1. Field of Invention
The present invention relates to the composition design and processing methods of the FeMnAlC alloys; and particularly to the methods of fabricating FeMnAlC alloys which simultaneously exhibit high strength, high ductility, and high corrosion resistance.
2. Description of the Prior Art
Austenitic FeMnAlC alloys have been subjected to extensive researches over the last several decades, because of their promising application potential associated with the high mechanical strength and high ductility. In the FeMnAlC alloy systems, both Mn and C are the austenite-stabilizing elements. The austenite (γ) phase has a face-center-cubic (FCC) structure; while Al is the stabilizer of the ferrite (α) phase having a body-center-cubic (BCC) structure. Hence, by properly adjusting the contents of the three alloying elements, it is possible to obtain fully austenitic FeMnAlC alloys at room temperature. Prior arts showed that the microstructure of the FeMnAlC alloys with a chemical composition in the range of Fe-(26-34) wt. % Mn-(6-11) wt. % Al-(0.54-1.3) wt. % C was purely single γ-phase without any precipitates after the alloys were solution heat-treated at 980-1200° C. and then quenched to room-temperature or ice water. Depending on the chemical composition, the ultimate tensile strength (UTS), yield strength (YS), and elongation of the as-quenched alloys were 814˜993 MPa, 423˜552 MPa, and 72-50%, respectively. These results indicate that, although it is possible to obtain single γ-phase with excellent ductility in as-quenched FeMnAlC alloys by properly adjusting the alloy compositions, the mechanical strength of these alloys is relatively low. Thus, prior arts are unable to achieve the goal of obtaining alloys that simultaneously possess high mechanical strength and high ductility in the as-quenched state.
In order to improve the mechanical strength of the Fe—Mn—Al—C alloys, prior arts have revealed that when the as-quenched alloys were aged at 500-650° C. for moderate times, a high density of fine (Fe,Mn)3AlCx carbides (so-called κ′-carbides) was found to precipitate coherently within the austenite matrix. The κ′-carbide has an ordered face-center-cubic (FCC) L′12 crystal structure. From these extensive studies disclosed in the prior arts, the significant improvement of the mechanical strength obtained in the aged FeMnAlC alloys is mainly due to the coherent precipitation of the fine κ′-carbides. However, since the κ′-carbides are rich in carbon and aluminum, the precipitation of these carbides from the supersaturated austenite matrix involves diffusion process of large amount of carbon and relevant alloy elements. Consequently, longer aging time and/or higher aging temperature are usually required. From numerous studies reported previously, an optimal combination of strength and ductility for the FeMnAlC alloys could be obtained through aging treatment at 550° C. for 15˜16 hours. This is primarily because that under these treatment conditions, a tremendous amount of fine κ′-carbides was found to precipitate within the austenite matrix and no precipitates were formed on the grain boundaries. According to the prior arts, depending on the alloy compositions, the UTS, YS and El of the FeMnAlC alloys aged at 550° C. for 15˜16 hours can reach 1130˜1220 MPa, 890˜1080 MPa and 39˜31.5%, respectively. However, if the aging process was performed at 450° C., it may take more than 500 hours to reach the same level of mechanical strength. Similarly, for 500° C. aging treatment, 50˜100 hours were needed.
In another embodiment, prior arts also tried to prolong the aging time at 550˜650° C. However, it was found that prolonged aging not only resulted in the growth of the fine κ′-carbides but also led to the γ→γ0+κ, γ0+κ, γ→α+κ, γ→κ+β-Mn, or γ→α+κ+β-Mn reactions occurring on grain boundaries. Where γ0 is the carbon-depleted γ phase and the κ-carbides have the same ordered FCC L′12 structure as the κ′-carbide, except that they usually precipitate on the grain boundaries with larger size. [Note: Conventionally, for distinction purpose, the finer (Fe,Mn)3AlCx carbides formed within the austenite matrix are termed as “κ′-carbides”, while the coarser (Fe,Mn)3AlCx carbides formed on the grain boundaries are termed as “κ-carbides”.] As a result, prolonged aging treatments frequently resulted in embrittlement of the alloys due to the precipitation of coarse κ-carbides on the grain boundaries.
The following publications gave more detailed descriptions and discussions of the abovementioned characteristics [1]-[20].    (1) S. M. Zhu and S. C. Tjong: Metall. Mater. Trans. A. 29 (1998) 299-306. (2) J. S. Chou and C. G. Chao: Scr. Metall. 26 (1992) 261-266. (3) T. F. Liu, J. S. Chou, and C. C. Wu: Metall. Trans. A. 21 (1990) 1891-1899. (4) S. C. Tjong and S. M. Zhu: Mater. Trans. 38 (1997) 112-118. (5) S. C. Chang, Y. H. Hsiau and M. T. Jahn: J. Mater. Sci. 24 (1989) 1117-1120. (6) K. S. Chan, L. H. Chen and T. S. Liu: Mater. Trans. 38 (1997) 420-426. (7) J. D. Yoo, S. W. Hwang and K. T. Park: Mater. Sci. Eng. A. 508 (2009) 234-240. (8) H. J. Lai and C. M. Wan: J. Mater. Sci. 24 (1989) 2449-2453. (9) J. E. Krzanowski: Metall. Trans. A. 19 (1988) 1873-1876. (10) K. Sato, K. Tagawa and Y. Inoue: Scr. Metall. 22 (1988) 899-902. (11) K. Sato, K. Tagawa and Y. Inoue: Mater. Sci. Eng. A. 111 (1989) 45-50. (12) I. Kalashnikov, O. Acselrad, A. Shalkevich and L. C. Pereira: J. Mater. Eng. Perform. 9 (2000) 597-602. (13) W. K. Choo, J. H. Kim and J. C. Yoon: Acta Mater. 45 (1997) 4877-4885. (14) K. Sato, K. Tagawa and Y. Inoue: Metall. Trans. A. 21 (1990) 5-11. (15) S. C. Tjong and C. S. Wu: Mater. Sci. Eng. 80 (1986) 203-211. (16) C. N. Hwang, C. Y. Chao and T. F. Liu: Ser. Metall. 28 (1993) 263-268. (17) C. Y. Chao, C. N. Hwang and T. F. Liu: Scr. Metall. (1993) 109-114. (18) T. F. Liu and C. M. Wan, Strength Met. Alloys, 1 (1986) 423-427. (19) G. S. Krivonogov, M. F. Alekseyenko and G. G. Solov'yeva, Fiz. Metal. Metallov ed., 39, No. 4 (1975) 775-781. (20) R. K. You, P. W. Kao and D. Gran, Mater. Sci. Eng., A117 (1989) 141-147.
Another method disclosed in the prior arts to further enhance the strength was adding small amounts of V, Nb, W and Mo to the austenitic FeMnAlC (C≦1.3 wt. %) alloys. After solution heat-treatment or controlled-rolling followed by an optimal aging at 550° C. for about 16 hrs, the UTS, YS, and El of the Fe-(25-31) wt. % Mn-(6.3-10) wt. % Al-(0.6-1.75) wt. % M(M=V, Nb, W, Mo)-(0.65-1.1) wt. % C alloys were significantly increased up to 953˜4259 MPa, 910˜1094 MPa, and 41˜26%, respectively.
The following publications gave more detailed descriptions and discussions of the abovementioned characteristics [21]-[25].
    (21) I. S. Kalashnikov, B. S. Ermakov, O. Aksel'rad and L. K. Pereira, Metal. Sci. Heat. Treat. 43 (2001) 493-496. (22) I. S. Kalashnikov, O. Acselrad, A. Shalkevich, L. D. Chumakova and L. C. Pereira, J. Mater. Proc. Tech. 136 (2003) 72-79. (23) K. H. Han, Mater. Sci. Eng. A 279 (2000) 1-9. (24) G. S. Krivonogov, M. F. Alekseyenko and G. G. Solov'yeva, Fiz. Metall. Metalloved. 39 (1975) 775. (25) I. S. Kalashnikov, B. S. Ermakov, O. Aksel'rad and L. K. Pereira, Metal. Sci. Heat. Treat. 43 (2001) 493-496.
Obviously, the Fe-(28-34) wt. % Mn-(6-11) wt. % Al-(0.54-1.3) wt. % C and Fe-(25-31) wt. % Mn-(6.3-10) wt. % Al-(0.6-1.75) wt. % M (M=V, Nb, W, Mo)-(0.65-1.1) wt. % C alloys disclosed in the prior arts and published literature can possess excellent combinations of mechanical properties, namely high-strength and high-ductility. However, they generally exhibited poor corrosion resistance. For instance, for the abovementioned alloys, the corrosion potential (Ecorr) and pitting potential (Epp) in the 3.5% NaCl aqueous solution (mimicking the sea water environment) were within the ranges of Ecorr=−750˜−900 mV and Epp=−350˜−500 mV, respectively. This strongly indicates that the alloys do not have adequate corrosion resistance when serving in sea water environment. In order to enhance the corrosion resistance, previous studies had added Cr to the alloys. It was pointed out that, by adding 3-9 wt. % of Cr, the corrosion resistance of the alloys could be significantly improved and an apparent passivation region can be observed in the current-voltage polarization curves. Previous results indicated that, by adding more than 3.3 wt. % of Cr to the Fe-(28-34) wt. % Mn-(6.7-10.5) wt. % Al-(0.7-1.2) wt. % C alloys, a significant improvement in corrosion resistance could be obtained. For instance, previous studies on Fe-30 wt. % Mn-9 wt. % Al-(3,5,6.5,8) wt. % Cr-1 wt. % C alloys have revealed a remarkable improvement in alloy's corrosion resistance when the Cr concentration exceeded 3.5 wt. %. When the Cr concentration was up to 5 wt. %, the alloys under the as-quenched condition exhibited an improvement of Ecorr and Epp to −560 mV and −50 mV in 3.5% NaCl solution, respectively. However, when the Cr concentration was increased to 6.5 and 8.0 wt. %, the corrosion resistance of the alloys decreased with increasing Cr concentration: Ecorr=−601 mV and Epp=−308 mV for Cr=6.5 wt. %; Ecorr=−721 mV and Epp=−380 mV for Cr=8.0 wt. %, respectively. Additionally, in the previous study concerning the corrosion behaviors of the Fe-30 wt. % Mn-7 wt. % Al-(3,6,9) wt. % Cr-1.0 wt. % C alloys in 3.5% NaCl solution, it was reported that when the Cr concentration was increased to about 6 wt. o, the Ecorr and Epp of the as-quenched alloy could be improved to −556 mV and −27 mV, respectively. However, when the Cr concentration was increased to 9 wt. %, the Ecorr and Epp of the as-quenched alloy were dramatically decreased to −754 mV and −472 mV, respectively. Investigations disclosed in the prior arts have pointed out that the Cr≦6 wt. % addition could be completely dissolved in Fe-30 wt. % Mn-7 wt. % Al-1.0 wt. % C alloy at the solution heat-treatment temperature of 1100° C. Consequently, the corrosion resistance of the alloys could be pronouncedly improved with increasing Cr concentration. However, when the Cr concentration was increased up to 9 wt. %, the Cr-rich carbides could be detected in the as-quenched alloy. The formation of the Cr-rich carbides resulted in the drastic decrease of the Ecorr and Epp values. In particular, it should be emphasized here that, even under the optimal composition conditions giving rise to the best corrosion resistance, such as alloys with the composition of Fe-30 wt. % Mn-7.0 wt. % Al-6.0 wt. % Cr-1.0 wt. % C, its performance in corrosion resistance is still far below those of AISI 304 (in 3.5% NaCl solution Ecorr=−350˜−210 mV, Epp=+100˜+500 mV) and AISI 316 (Ecorr=−200 mV, Epp=+400 mV) austenitic stainless steels or the 17-4PH precipitation-hardening stainless steels (Ecorr=−400˜−200 mV, Epp=+40˜+160 mV).
Moreover, since Cr is a very strong carbide former, prior arts have shown that, although the as-quenched alloys usually reveal single austenite phase when the Cr concentration is below about 6 wt. %, coarse Cr-rich carbides, such as (Fe,Mn,Cr)23C6 and (Fe,Mn,Cr)7C3, can easily precipitate on the grain boundaries during the aging treatment. As a result, the aged alloys frequently exhibit dramatic reduction in both their ductility and corrosion resistance. This is also the primary reason why most of the austenitic Fe—Mn—Al—Cr—C alloys disclosed in the prior arts or published literature have been used in the as-quenched condition and seldom carried out any aging treatment. In a series of Fe-(26.5-30.2) wt. % Mn-(6.85-7.53) wt. % Al-(3.15-9.56) wt. % Cr-(0.69-0.79) wt. % C alloys disclosed in the prior arts, the UTS and YS of the alloys are respectively ranging within 723˜986 MPa and 410˜635 MPa after solution heat-treatment. If one compares these mechanical properties with those of the abovementioned Fe—Mn—Al—C alloys subjected to 15˜16 hours of aging at 550° C. (UTS=1130˜1220 MPa YS-890˜1080 MPa), it is apparent that, although exhibiting superior corrosion resistance, the austenitic Fe—Mn—Al—Cr—C alloys have much lower mechanical strength than the aged Fe—Mn—Al—C alloys.
The following publications gave more detailed descriptions and discussions of the abovementioned characteristics [26]-[39].    (26) C. Y. Chao, 2001, “Low density high ductility Fe-based alloy materials for golf club heads”, Patent No. 460591, Taiwan, R.O.C. (27) C. Y. Chao, 2004, “Low density Fe-based materials for golf club heads”, Patent No. 460591, Taiwan, R.O.C. (Same as US Patent No.: US006007). (28) T. F. Liu and J. W. Lee, 2007, “Low density, high strength, high toughness alloy materials and the methods of making the same”, Patent No. I279448, Taiwan, R.O.C. (29) Tai W. Kim, Jae K. Han, Rae W. Chang and Young G. Kim, 1995, “Manufacturing process for austenitic high manganese steel having superior formability, strengths and weldability”, U.S. Pat. No. 5,431,753. (30) C. S. Wang, C. Y. Tsai, C. G. Chao and T. F. Liu: Mater. Trans. 48 (2007) 2973-2977. (31) S. C. Chang, J. Y. Liu and H. K. Juang: Corros. Eng. 51 (1995) 399-406. (32) S. C. Chang, W. H. Weng, H. C. Chen, S. J. Liu and P. C. K. Chung: Wear 181-183 (1995) 511-515. (33) C. J. Wang and Y. C. Chang: Mat. Chem. Phy. 76 (2002) 151-161. (34) J. B. Duh, W. T. Tsai and J. T. Lee, Corrosion November (1988) 810. (35) M. Ruscak and T. R. Perng, Corrosion 51 (1995) 738-743. (36) C. J. Wang and Y. C. Chang, Mater. Chem. Phy. 76 (2002) 151-161. (37) S. T. Shih, C. Y. Tai and T. P. Perng, Corrosion February 49 (1993) 130-134. (38) Y. H. Tuan, C. S. Wang, C. Y. Tsai, C. G. Chao and T. F. Liu: Mater. Chem. Phy. 114 (2009) 246-249. (39) Y. H. Than, C. L. Lin, C. G. Chao and T. F. Liu: Mater. Trans. 49 (2008) 1589-1593.
The characteristics of the Fe-(26-34) wt. % Mn-(6-11) wt. % Al-(0.54-1.3) wt. % C and Fe-(25-31) wt. % Mn-(6.3-10) wt. % Al-(0.6-1.75) wt. % M(M=V,Nb,Mo,W)-(0.65-1.1) wt. % C alloys disclosed in the prior arts can be summarized as following. For alloys containing less than 1.4 wt. % of carbon, the microstructure of the alloys after being solution heat-treated at 980˜1200° C. and then quenched, is single austenite phase or austenite phase with small amount of (V, Nb)C carbides. When the as-quenched alloys are aged at 550° C. for 15˜16 hours, the alloys can achieve the optimal combination of high-strength and high-ductility. However, the alloys usually exhibit poor corrosion resistance. When up to approximately 6 wt. % of Cr was added to the austenitic Fe—Mn—Al—C alloys, the corrosion resistance can be improved in the as-quenched condition. Nevertheless, due to the precipitation of coarse Cr-rich carbides on the austenite grain boundaries during aging treatments, the alloys easily lose their ductility and corrosion resistance. Therefore, it can be concluded from the above discussions that the compositions of various Fe—Mn—Al—C, Fe—Mn—Al-M (M=V, Nb, W, Mo)—C, and Fe—Mn—Al—Cr—C alloys and the associated processing conditions disclosed in the prior arts have failed to accomplish the goal of producing an alloy possessing the characteristics of high-strength, high-ductility, and high corrosion resistance, simultaneously.
In order to overcome these unresolved outstanding problems, the present inventor, based on decades of practical experiences in materials researches, including alloy designs and technology developments of Fe—Mn—Al—C alloys, has carried out numerous of experiments and come up with the present novel invention.